Formation of Ultrafine Grained Dual Phase Steels through Rapid Heating

Transcription

1 , pp Formation of Ultrafine Grained Dual Phase Steels through Rapid Heating Hamid AZIZI-ALIZAMINI, Matthias MILITZER and Warren J. POOLE The Centre for Metallurgical Process Engineering, The University of British Columbia, Vancouver, BC, Canada V6T 1Z4. (Received on January 31, 2011; accepted on March 16, 2011) In this study, ultrafine grained dual phase structures have been developed in a plain low carbon steel, 0.17C and 0.74Mn (wt pct). The approach is based on rapid heating of a very fine ferrite-carbide aggregate into the intercritical annealing region followed by water quenching. This rapid heat treatment results in an ultrafine grained dual phase steel with improved properties. The effect of thermomechanical processing parameters such as heating rate and intercritical annealing time on the microstructure and mechanical properties have been examined. The key factors contributing to the grain refinement are uniform distribution of nanosize cementite particles acting as potential sites for austenite nucleation as well as the limited time available for coarsening of the microstructure. The mechanical properties of the present ultrafine grained dual phase steel show an excellent combination of strength and uniform elongation because of considerable work hardening. KEY WORDS: plain low carbon steels; rapid heating; ultrafine grain; intercritical annealing; dual phase. 1. Introduction Recent demands for lighter vehicles resulted in the development of advanced high strength steels (AHSS). The ultra light steel auto body-advanced vehicle concept (ULSAB- AVC) 1) and SuperLIGHT-CAR 2) studies are examples of the global interest in this subject. The usage of AHSSs including dual phase, complex phase and transformation induced plasticity (TRIP) steels is indeed gradually increasing and is expected to rise to 35 wt pct by ) Dual phase steels offer a unique combination of mechanical properties making them one of the materials of choice for sheet products in the automotive industry. Continuous yielding, high work hardening rate and relatively high formability are some of the essential properties for many automotive applications. However, further improvement in properties remains essential in the competitive automotive industry. One way to increase the strength in steels without additional alloying is grain refinement. However, this approach results in detrimental effects such as diminishing the work hardening rate and introducing Lüders bands in ferritic steels that hamper their potential application for structural purposes. The yield ratio for ultrafine grained (UFG) ferritic steels with average grain sizes between 1 3 μ m is usually above ) But, a different trend is reported for refining the microstructure in dual phase steels. Balliger and Gladman 10) have shown that refining martensite islands in dual phase steels results in an increase in work hardening rate. Tsipouridis et al. 11) revealed that the strength of dual phase steels can be improved via grain refinement while the ductility remains unaffected in a Mo Cr alloyed steel. The yield ratio of 0.59 for a C Mn UFG dual phase steel is comparable with the coarse grained counterparts. 12) Further, Park et al. 13) showed that the desired mechanical properties of conventional coarse grained (CG) dual phase (DP) steels such as continuous yielding and low yield ratio can be translated into the UFG scales. The potential for further improvement of the mechanical properties in dual phase steels via grain refinement has recently stimulated the investigations of techniques by which UFG dual phase steels can be achieved. Park et al. 13) utilized an equal channel angular pressed (ECAPed) ferritepearlite microstructure as a starting structure before intercritical annealing to produce UFG dual phase steels. Yet, this processing approach remains practically challenging for sheet products. Calcagnotto et al. 12) showed that intercritical annealing of a low carbon steel with 0.17C and 1.6Mn (wt pct) processed via large strain warm deformation can also lead to a UFG dual phase structure with improved properties. However, applying the same technique to a leaner steel chemistry, i.e. with 0.87 wt pct Mn, resulted in a relatively CG structure consisting of ferrite, martensite and pearlite. Mukherjee et al. 14) successfully developed UFG dual phase structures through the industrially more suitable technique of strain-induced transformation. However, this approach would be applicable to hot-rolled products rather than coldrolled and coated sheets that are of primary interest to the current work. Rapid heating and cooling experiments have also been used to develop very fine structures in steels using rapid transformation annealing (RTA) 15) and flash processing. 16) Andrade-Carozzo and Jacques 17) showed that rapid heating of an initial cold-rolled ferrite-pearlite structure into the intercritical annealing region results in grain refinement in the final dual phase structures of a Nb-bearing low carbon steel. However, there have been a very limited number of 2011 ISIJ 958

2 studies to develop UFG dual phase steels in plain low carbon steels ,18 20) Thus, the present work seeks to examine combinations of cold rolling and annealing processing steps to develop UFG dual phase structures in a plain low carbon steel with 0.17C and 0.74Mn (wt pct). A systematic study is presented to investigate the effect of initial structure and processing parameters such as heating rate and intercritical annealing time on the microstructure evolution and resulting mechanical properties. 2. Materials and Experimental Procedure A hot-rolled low carbon steel with the chemical composition given in Table 1 was used in this study. Figure 1 shows representative micrographs of the as-received hotrolled structure that consists of 80.0±1.7 vol pct ferrite and 20.0±1.7 vol pct pearlite. Three different thermomechanical processes were designed to produce different microstructures, i.e. I) cold-rolled martensite, II) cold-rolled tempered martensite and III) fine grained ferrite-carbide aggregate. In each case, the final processing step involves rapid heating to the intercritical region followed by water quenching. The difference is related to the processing prior to intercritical annealing. Figures 2(a) 2(c) show these processes schematically. These processing steps were chosen to vary the initial structure systematically prior to intercritical annealing. In process (I), Fig. 2(a), the initial ferrite-pearlite structure was first austenitized at C for 30 min followed by ice brine quenching to assure a fully martensitic structure. It was then Table 1. Chemical composition of steel used in this study. Elements Fe C Mn P S Si Al N wt pct Bal Fig. 1. (a) Optical and (b) SEM micrographs of initial hot-rolled ferrite-pearlite structure (F: ferrite; P: pearlite; C: carbide). 80 pct cold rolled with a laboratory rolling mill (roll diameter: 130 mm). Process (II) is similar to process (I), Fig. 2(b), however, the martensite structure was tempered at 550 C for 2 h prior to 80 pct cold reduction. In process (III), Fig. 2(c), after annealing the martensite structure at 550 C for 1 h followed by 80 pct cold rolling, an additional 75 min annealing at 550 C was performed. All annealing treatments were conducted in a tube furnace with Ar controlled atmosphere. The samples obtained from those different processes were then rapidly heated into the intercritical annealing region at 300 C/s and held for 10 s at 750 C followed by water quenching. Through analyzing and comparing the microstructures and mechanical properties of the resulting dual phase structures, the best processing route was then selected for a detailed evaluation of the intercritical annealing conditions. The effect of intercritical processing parameters such as heating rate and holding time on microstructure and properties were systematically investigated. All intercritical annealing treatments were done at 750 C. Different heating rates of 1 C/s, 50 C/s and 300 C/s were used for comparison purposes and the holding time was varied in the range of s. Test coupons of mm were machined from the cold-rolled sheets with the longitudinal direction of the test coupon being aligned with the rolling direction. All controlled heating rate experiments were conducted in a Gleeble 3500 thermomechanical simulator. The testing chamber was first put under vacuum, 0.04 Pa ( Torr), and then back filled with Ar. The temperature was controlled using a type K thermocouple spot-welded on the center of the sample. A mechanical dilatometer was attached to the center of the samples to measure the change in width during heating. The volume fraction of austenite was determined via analyzing the dilatometric data employing the lever rule, details of which can be found elsewhere. 21) All microstructures were analyzed in the transverse plane of the rolled sheet, i.e. the through thickness plane perpendicular to the rolling direction. Microstructural characterization was carried out using optical and electron microscopy using a Hitachi S2300 scanning electron microscope (SEM) with a secondary electron detector and a 200 kv Hitachi H- 800 transmission electron microscope (TEM). Electron backscatter diffraction (EBSD) analysis was performed using a Schottkey source field emission scanning electron microscope, FEI XL30 SEM. The EBSD data analysis was performed with HKL Channel 5 software. For SEM and Fig. 2. Thermomechanical processes employed to develop different initial structures (IBQ: Ice Brine Quench, WQ: Water Quench, CR: Cold Rolling) ISIJ

3 EBSD analyses, samples were electropolished in 95 pct acetic acid and 5 pct perchloric acid solution. Etching with 3 pct Nital enabled to reveal the microstructure for optical and conventional SEM samples. For TEM observation, thin foils were prepared by twin-jet polishing technique using a mixture of 95 pct acetic acid and 5 pct perchloric acid at an applied potential of 40 V at 20 C. Measurements of grain size and volume fraction of ferrite and martensite phases were done using SEM micrographs. Grain size measurements were based on the equivalent area diameter (EQAD) approach and at least 500 grains were analyzed. The quantitative measurements were carried out using Clemex image analysis software. Further, to quantify the resulting mechanical properties, tensile tests were conducted on samples with 12.5 mm gauge length at a nominal strain rate of s 1 using a MTS servo-hydraulic machine. 3. Results 3.1. Microstructure Evolution Initial Structures The common feature of all heat treatments is the initial martensite structure. Martensite is intrinsically a fine scale structure with several levels of hierarchy in its microstructure. 8) Further, its highly dislocated fine structure with supersaturated carbon content provides all the necessary ingredients for further microstructure evolution such as precipitation and recrystallization. Therefore all the designed scenarios for developing initial structures depicted in Fig. 2 start with a fully martensitic structure. Figure 3(a) shows an optical micrograph of the as-quenched fully martensitic structure. TEM observation confirmed a lath martensite structure with a lath width of μ m, Fig. 3(b). A dislocation density of approximately m 2 has been reported for 0.18C (wt pct) steels with lath martensitic structure. 22) Figure 3(c) shows an EBSD orientation map of the asquenched martensite. High angle boundaries (> 15º misorientation) are indicated by black lines. The EBSD map clearly reveals the presence of fine features of lath martensite structure with different orientations (colors). Using the approach proposed by Ueji et al., 8) the mean grain size of lath martensite considering the boundaries with misorientation larger than 15 was measured. The representative grain size of martensite was 4.3 μm, and is comparable with the reported 3.2 μ m and 2.1 μm grain size for lath martensite in ) and ) carbon (wt pct) steels, respectively. The martensite structure was then subjected to different thermomechanical procedures to create three different initial structures prior to intercritical annealing as follow: (I) Deformed martensite: Figure 4(a) shows an 80 pct cold-rolled martensite structure as initial structure for process (I). Deformation features such as kinked lath structures and bent lamellae 8) can be seen in this micrograph. Interstitial carbon atoms can be expected to mainly reside at the lath boundaries as well as dislocations. 24) (II) Deformed ferrite-carbide aggregate: Tempering prior to cold rolling resulted in the presence of nano-size carbide particles in the ferrite matrix that can be seen in Fig. 4(b). This microstructure is the initial structure for process (II). The major difference with the previous initial structure, i.e. deformed martensite, is the presence of carbide particles. In this scenario, the deformed ferrite matrix is elongated in the rolling direction as can be seen in Fig. 4(b). (III) Equiaxed ferrite-carbide aggregate: Further annealing of the deformed ferrite-carbide structure at 550 C for 75 min resulted in a UFG ferrite-carbide aggregate with equiaxed ferrite grains (Fig. 4(c)) which is the initial structure for process (III). The annealing time was optimized via a series of isothermal annealing experiments. A lower dislocation density and the equiaxed morphology of ferrite grains are the major differences between this structure and the deformed ferrite-carbide aggregate presented in Fig. 4(b). The inset in Fig. 4(c) shows the bimodal nature of size distribution of nano-size carbide particles. Larger particles with an approximate average size of ~150 nm were mainly located at ferrite grain boundaries while smaller particles with an average size of ~50 nm were distributed inside the grains. Fig. 3. (a) optical micrograph of as-quenched martensite structure, (b) bright field (BF) TEM image of lath martensite structure with the corresponding selected area diffraction (SAD) pattern and (c) EBSD orientation map of as-quenched martensite together with superimposed boundary map. Fig. 4. Initial microstructures developed for intercritical annealing treatment: (a) deformed martensite structure (process I), (b) deformed ferrite-carbide aggregate (process II) and (c) UFG ferrite-carbide aggregate (process III) ISIJ 960

4 Benchmarking for Selection of Initial Structure Samples with the three microstructures shown in Fig. 4 were then heated at 300 C/s to 750 C and held for 10 s followed by water quenching. Starting with initial structures obtained from the processes (I) and (II), the typical resulting dual phase microstructures can be seen in Figs. 5(a) and 5(b), respectively. These microstructures consist of two distinct regions: UFG dual phase structure with uniform distribution of ferrite and martensite phases and relatively coarse ferrite grains elongated in the rolling direction. However, a homogeneous UFG dual phase microstructure was produced from the UFG ferrite-carbide aggregate developed in process (III), Fig. 5(c). Table 2 summarizes these characteristics of the microstructures. It is evident that at comparable volume fraction of martensite, average sizes for ferrite grains and martensite islands are approximately equal. Since the volume of the individual grains is of primary importance when considering mechanical and fracture behaviour, the volumetric size distribution of the ferrite grains should be considered as shown in Fig. 6. Comparison between different distributions reveals an interesting observation. Assuming a ferrite grain size of 3 μm as a boundary between fine and coarse grain size regimes 25), then the volume fraction of coarse grains in the microstructures developed in processes (I) and (II) is 0.7 and 0.6, respectively. But, this value is reduced to 0.32 in the UFG dual phase structure developed in process (III). These distributions strongly influence the corresponding mechanical properties including the fracture behavior as presented in Table 3. A better combination of strength and ductility is achieved in the DP structure obtained from process (III), i.e. when predominantly a truly UFG dual phase structure is attained. The improvement can be clearly highlighted in the fracture behavior since strain to fracture is more than doubled in the UFG dual phase structure compared to the coarser dual phase microstructure developed in process (I). As a result, the UFG ferrite-carbide aggregate obtained from process (III) is selected as the initial structure for further investigations Detailed Examination of the UFG DP Structure Figure 7 shows bright field TEM micrographs of the UFG dual phase structure intercritically annealed at 750 C for 10 s with the corresponding diffraction pattern. Closer observation using higher magnification reveals dislocations in the ferrite grains especially adjacent to the martensite Fig. 6. Estimated volumetric ferrite grain size distribution for the microstructures obtained in (a) process (I), (b) process (II) and (c) process (III). The volume of each grain with a 2D EQAD, d, is estimated using V d = nd 3 /6. Fig. 5. Table 2. Microstructural evolution after 10 s annealing at 750 C for initial structures obtained from (a) process (I), (b) process (II) and (c) process (III). Comparison of the effect of initial microstructure on the final dual phase structure. Process Process (I) Process (II) Process (III) Martensite volume fraction Martensite island size (EQAD), μm Average ferrite grain size (EQAD), μm Table 3. Process Comparison of mechanical properties for different dual phase structures. YS (1), MPa (0.2 pct offset) UTS (2), MPa U. El. (3), (pct) **Fracture strain, ε f Process (I) Process (II) Process (III) (1) YS: yield strength, (2) UTS: ultimate tensile strength, (3) U.El.: uniform elongation. ** Fracture strain, ε f, was measured using ε f = ln (A 0/A f) where A 0 and A f are the initial cross section area and area of fracture surface of the tensile test samples, respectively. The fracture area, A f, was measured using SEM observation ISIJ

5 islands, Fig. 7(b), in agreement with observations made by Son et al. 26) and Rigsbee et al.. 27) For further examination, EBSD analysis was performed. Figure 8 shows the image quality (IQ) and grain boundary misorientation maps. Martensite islands can be distinguished from ferrite grains by their lower image quality and band contrast. The detailed procedure to distinguish ferrite from martensite is given by Mukherjee et al. 14) Following the proposed procedure, a critical band contrast of 55 (BC critical) was chosen to separate ferrite from martensite and to match the volume fraction of martensite measured from the SEM micrographs, i.e. 0.32, as presented in Table 2. The objects with values below this boundary are shown as black regions in Fig. 8(b). The grain boundary misorientation distribution was then analyzed for the ferrite regions. As shown in Fig. 8(d) more than 85 pct of the ferrite grain boundaries are high angle boundaries with misorientation above 15. The average size of these ferrite grains is approximately 2 μ m which is about twice the size obtained using SEM measurements. A similar discrepancy between the EBSD and the SEM grain size measurements in low carbon steels was also reported by Mukherjee et al. 14) This is essentially related to the presence of subgrains with misorientation angles Fig. 7. Bright field TEM images of UFG dual phase structure obtained by heating at 300 C/s and intercritically annealed at 750 C for 10 s followed by water quenching (M: martensite, F: ferrite). between 2 and 15, see Fig. 8(b). These subgrains can be mistakenly interpreted as grains in the SEM observations. Adding grain boundaries with 2º 15º misorientation to the grain size measurement using the EBSD analysis resulted in an average grain size of approximately 1.5 μm that is similar to that obtained in the conventional SEM measurements. 4. Discussion Results presented in the previous sections revealed that a fine aggregate of equiaxed ferrite grains with carbide particles is a starting microstructure of choice to develop UFG dual phase structures. However, there are two aspects of the proposed thermomechanical treatment that should be addressed. First, the role of heating rate need to be further investigated to see the possibility of employing more industrially applicable heating rates. Secondly, the effect of holding time on microstructure evolution should be determined to evaluate the processing window for the proposed intercritical annealing stage. Figure 9 presents the effect of heating rate on microstructure evolution. Two different sets of samples were heated to 750 C at 1 C/s and 50 C/s and then held for 10 s before water quenching. Figure 9(a) shows that at the lower heating rate, i.e. 1 C/s, substantial microstructure coarsening takes place due to the increased processing time. The volume fraction of martensite is 0.18 in this case. Average ferrite grain and martensite island sizes are 8 μ m and 5 μm, respectively. At 50 C/s, localized growth of ferrite grains results in a relatively coarse grained structure with patches of UFG ferrite-martensite structure, Fig. 9(b). Here, the volume fraction of martensite is Comparison of the microstructures achieved at different heating rates of 1 C/s, 50 C/s and 300 C/s, Figs. 5(c) and 9, shows that an increase in heating rate results in a rise in volume fraction of austenite from 0.18 to 0.32 for the given holding time of 10 s at 750 C. This trend can be explained when considering the effect of heating rate on ferrite grain growth and the associated density of austenite nuclei. Decreasing the heating rate promotes ferrite grain coarsening, see Fig. 9(a). It appears that austenite nucleates primarily at cementite particles located at ferrite grain boundaries. Further growth of austenite grains takes place along ferrite grain boundaries. Cementite particles inside the grains appear not to be viable nucleation sites at this heating rate. At higher heating rates, however, ferrite grain growth is greatly reduced. As a result, an increase in nucleation site density of austenite grains from m 2 to m 2 is estimated as the heating rate increases from 1 C/s to 300 C/s (this estimation is based on Fig. 8. EBSD results for UFG dual phase structure obtained by heating at 300 C/s followed by intercritical annealing at 750 C for 10 s, (a) Image quality (IQ) map, (b) phase map (ferrite: white, martensite: black) and (c) band contrast distribution, (d) ferrite grain boundary misorientation distribution with superimposed random distribution. Fig. 9. Microstructure evolution of intercritically annealed UFG ferrite-carbide aggregate at 750 C for 10 s with (a) 1 C/s and (b) 50 C/s heating rate followed by water quenching ISIJ 962

6 the assumption that each martensite island is an austenite nucleus turning into a grain with no coarsening effect). The morphology of austenite grains also changes to a more equiaxed one as the heating rate increases. Unlike at 1 C/s, there are very few carbide particles left undissolved at 300 C/s, compare Figs. 5(c) and 9(a). Figures 10(a) 10(c) show the effect of holding times of 1 s, 60 s and 300 s on the microstructure features at 750 C, respectively. All samples were heated at 300 C/s. With increasing the holding time substructures inside ferrite grains are removed and cementite particles, shown by arrows in Fig. 10(a), are mostly dissolved at the expense of austenite formation (martensite at room temperature). As summarized in Fig. 10(d), the volume fraction of martensite increases with holding time and reaches 0.35 after 60 s the paraequilibrium (PE) value for austenite (note: paraequilibrium denotes a constraint equilibrium without partitioning of substitutional alloying elements). When increasing the annealing time from 10 to 60 s, the size of martensite islands rises from approximately 1 μ m to 2 μm and the volume fraction of ferrite grains larger than 3 μm changes from 32 pct to 38 pct and reaches 46 pct after 300 s intercritical annealing. These findings reveal that high heating rate and short holding time are essential to obtain uniform UFG dual phase structures. Increasing the heating rate, equivalent to shorter ramping time, stimulates grain refinement in different ways as described below: At 1 C/s, substantial growth of ferrite grains upon heating results in coarser structure while at higher heating rates, i.e. 300 C/s, microstructural features of a UFG ferrite-carbide aggregate are mainly preserved prior to intercritical annealing. These features consist of high density of ferrite grain boundaries and the presence of a large population of carbide particles on these boundaries as shown in Fig. 4(c), both of which can be nucleation sites for austenite formation. Austenite preferentially nucleates at the cementite particles located at ferrite grain boundaries. 28) Judd and Paxton 29) showed that the nucleation rate can be three to eight times faster at ferrite grain boundaries compared to ferrite matrix. Thus, a combination of large population of ferrite grain boundaries and carbide particles can potentially increase the nucleation site density. Further, an increase in heating rate from 1ºC/s to 300ºC/s raises the superheat from 38ºC to 51ºC for measurable start of austenite formation, i.e. Ac 1- Ae 1 where the Ac 1 temperature is the austenite formation start temperature upon heating, measured using the dilatometry data, and the Ae 1 temperature is the equilibrium austenite formation start temperature (calculated using the Thermocalc data, 704ºC here). This increase in superheat for austenite formation may further promote an increased nucleation density for austenite grains. Kaluba et al. 30) have also shown that for sufficiently high superheat in rapid heating experiments, austenite formation is also viable at ferrite grain boundaries with no carbides. However, unlike for austenite decomposition upon cooling, during heating both the driving force for austenite formation and the carbon diffusion rate increase such that growth of austenite grains is accelerated. This combination accelerates nucleation and growth of austenite grains. The combination of these factors on nucleation and growth of austenite grains is consistent with the observed increase in volume fraction of austenite as the heating rate rises and this is in agreement with the findings by Andrade- Carozzo and Jacques. 17) Coarsening of the microstructure during holding at 750 C is a consequence of both the curvature effect and diffusion distances. The coarsening of the austenite grains is promoted by the comparatively short diffusion distances for carbon between individual grains and can be thought of as coarsening of carbon-rich precipitates. 31) It has been suggested that austenite nuclei formed during intercritical annealing can stabilize the microstructure and impede ferrite grain growth. 17) Thus, ferrite grain coarsening is synchronized with the coarsening of the austenite grains, i.e. relatively limited growth of ferrite grains has taken place even after 300 s holding at 750 C, Fig. 10, when compared to ferrite grain growth during slow heating before austenite forms, Fig. 9(a). Table 4 summarizes mechanical properties for different dual phase structures. It can be seen that with increasing the heating rate from 50 C/s to 300 C/s, both the YS and the UTS increase at comparable ductility and yield ratio. At the same time, comparable mechanical properties are achieved while increasing the holding time from 10 s to 30 s. The balance between the uniform elongation and the UTS for different dual phase steels is summarized in Fig. 11. The open symbols represent conventional CG dual phase steels and the closed ones are available data on UFG dual phase Table 4. Summary of mechanical properties for different dual phase structures. Fig. 10. Microstructure evolution of intercritically annealed UFG dual phase structure at 750 C for (a) 1 s, (b) 60 s and (c) 300 s holding times followed by water quenching, (d) change in volume fraction of martensite and coarse ferrite grains at 750 C with paraequilibrium (PE) prediction of austenite fraction (M: martensite, F: ferrite). Heating rate, holding time ( C/s, sec) Martensite fraction U. El. (pct) YS, MPa (0.2 pct offset) UTS, MPa Yield ratio 50, , , , , ISIJ

7 industries to enable vehicle weight reduction while maintaining the performance intact. Acknowlegment The authors would like to acknowledge the financial support of the Natural Sciences and Engineering Research Council of Canada (NSERC). We are grateful to D. Field for valuable assistance in the EBSD measurements. REFERENCES Fig. 11. steels. 12,13,18,26,27,32 35) The results for UFG dual phase structures in this study are comparable to those reported in the literature. These data are at or above the upper limit for the balance in CG dual phase steels depicted by the solid line. UFG dual phase steels offer the opportunity to increase tensile strength by up to 150 MPa at comparable ductility, promising a new boundary that is shown by the dotted line. Available data on UFG dual phase steels is so far restricted to the upper tail of the balance, i.e. elongation values below 15 pct since these studies were carried out on steels with higher carbon content ( wt pct C) with relatively large volume fraction of martensite. Carbon primarily controls the strength of martensite. Using lower carbon content UFG dual phase steels or reducing the martensite fraction may provide data for UFG dual phase steels with higher uniform elongation values. 5. Summary Balance between UTS and uniform elongation for conventional CG 27,32 35) and UFG 12,13,18,19,26) dual phase steels. This study aimed to produce and characterize UFG dual phase structures in a plain low carbon steel with 0.17 wt pct carbon and 0.74 wt pct manganese. The proposed thermomechanical technique is based on rapid heating and cooling of a UFG ferrite-carbide aggregate that is originally developed from an initial martensite structure. It was shown that a rapid heating and cooling cycle is essential to guarantee UFG dual phase structures with optimum properties. This necessity is mainly to overcome undesired growth of ferrite and austenite grains that otherwise result in a coarser ferritemartensite structure. Short intercritical annealing treatment can be an advantage of this rather simple technique; however industrial implementation can be challenging with the existing production lines for steel sheets. One way to tackle the problem is to add appropriate alloying addition such as Cr, Mo and Nb to stabilize an initial fine grained structure and to impede microstructure coarsening during the heat treatment cycle. Assessment of the mechanical properties reveals that UFG dual phase steels have the capacity to replace conventional CG dual phase steels. The main beneficial features of dual phase steels can be translated into finer scale structures. 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